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Aligning graphene in bulk copper : nacre-inspired nanolaminated architecture coupled with in-situ processing for enhanced mechanical properties and high electrical conductivity
Reference: Cao Mu, Xiong Ding-Bang, Tan Zhanqiu, Ji Gang, Amin-Ahmadi Behnam, Guo Qiang, Fan Genlian, Guo Cuiping, Li Zhiqiang, Zhang Di.- Aligning graphene in bulk copper : nacre-inspired nanolaminated architecture coupled w ith in-situ processing for enhanced mechanical properties and high electrical conductivity Carbon - ISSN 0008-6223 - 117(2017), p. 65-74 Full text (Publisher's DOI): https://doi.org/10.1016/J.CARBON.2017.02.089
Institutional repository IRUA Accepted Manuscript
Aligning graphene in bulk copper: Nacre-inspired nanolaminated architecture coupled with in-situ processing for enhanced mechanical properties and high electrical conductivity
Mu Cao, Ding-Bang Xiong, Zhanqiu Tan, Gang Ji, Behnam Amin-Ahmadi, Qiang Guo, Genlian Fan, Cuiping Guo, Zhiqiang Li, Di Zhang PII: S0008-6223(17)30219-1 DOI: 10.1016/j.carbon.2017.02.089 Reference: CARBON 11793
To appear in: Carbon
Received Date: 29 December 2016 Revised Date: 22 February 2017 Accepted Date: 25 February 2017
Please cite this article as: M. Cao, D.-B. Xiong, Z. Tan, G. Ji, B. Amin-Ahmadi, Q. Guo, G. Fan, C. Guo, Z. Li, D. Zhang, Aligning graphene in bulk copper: Nacre-inspired nanolaminated architecture coupled with in-situ processing for enhanced mechanical properties and high electrical conductivity, Carbon (2017), doi: 10.1016/j.carbon.2017.02.089.
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Graphical Abstract
MANUSCRIPT
ACCEPTED ACCEPTED MANUSCRIPT
Aligning graphene in bulk copper: nacre-inspired nanolaminated architecture coupled with in-situ processing for enhanced mechanical properties and high electrical conductivity
Mu Cao a, Ding-Bang Xiong a, *, Zhanqiu Tan a, Gang Ji b, Behnam Amin-Ahmadi c, Qiang Guo a, Genlian Fan a, Cuiping Guo a, Zhiqiang Li a, and Di Zhang a, ** a State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai
200240, China MANUSCRIPT b Unité Matériaux et Transformations (UMET) CNRS UMR 8207, Université Lille1, 59655
Villeneuve d’Ascq, France c Electron Microscopy for Materials Science (EMAT), University of Antwerp,
Groenenborgerlaan 171, 2020- Antwerp, Belgium
* Corresponding author. E-mail address : [email protected] (Ding-Bang Xiong)
** Corresponding author. E-mail address : [email protected] (Di Zhang) ACCEPTED
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Abstract:
Methods used to strengthen metals generally also cause a pronounced decrease in ductility and electrical conductivity. In this work a bioinspired strategy is applied to surmount the dilemma.
By assembling copper submicron flakes cladded with in-situ grown graphene, graphene/copper matrix composites with a nanolaminated architecture inspired by a natural nacre have been prepared. Owing to a combined effect from the bioinspired nanolaminated architecture and improved interfacial bonding, a synergy has been achieved between mechanical strength and ductility as well as electrical conductivity in the graphene/copper matrix composites. With a low volume fraction of only 2.5% of graphene, the composite shows a yield strength and elastic modulus ~177% and ~25% higher than that of unreinforced copper matrix, respectively, while retains ductility and electrical conductivity comparable to that of pure copper. The bioinspired nanolaminated architecture enhances the efficiencieMANUSCRIPTs of two-dimensional (2D) graphene in mechanical strengthening and electrical conducting by aligning graphene to maximize performance for required loading and carrier transporting conditions, and toughens the composites by crack deflection. Meanwhile, in-situ growth of graphene is beneficial for improving interfacial bonding and structural quality of graphene. The strategy sheds light on the development of composites with good combined structural and functional properties.
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1. Introduction
Analogous to that transforming crystal structure at atomic and molecular scale can lead to change in the properties of a compound, recently tailoring properties by architecture design that changes the spatial distribution of reinforcement in matrix at micro-/nano-scale without changing constituents has attracted intensive attention in the field of composites [1]. Natural biological materials are usually made up of only simple constituents, but show remarkable range of mechanical and functional properties, which can be attributed to multiscale transformation on their architectures [2-6]. Understanding the role that multilevel architectures play in controlling properties of natural materials may serve as inspirations for architecture design in composites.
Usually, conventional metal matrix composites (MMCs) containing homogeneously distributed reinforcement exhibit higher strength but compromise ductility and toughness compared to the pure matrix. In nature, the most MANUSCRIPTtypical example that surmounts the conflict between strength and toughness by architecture design is probably the nacreous part of seashells
[7]. Nacre is made up of about 95 vol % brittle mineral aragonite (a polymorph of calcium
carbonate) and only a few percent of soft organic material, however, it exhibits phenomenal
fracture strength and toughness properties thanks to a perfectly ordered “brick-and-mortar”
architecture [8]. For example, its work of fracture is three orders of magnitude greater than that of a single crystal of its constituent mineral aragonite [9]. Therefore, mimicking the unique
“brick-and-mortar” architecture in nacre might be a promising strategy for producing MMCs with optimum combinationACCEPTED of strength and ductility and toughness. Such bioinspired architecture could be built in MMCs by using high strength reinforcement with high aspect ratio as “brick” combined with ductile metal “mortar”. While “brick” of reinforcement is by no means easy to obtain, some emerging two-dimensional (2D) nanomaterials show great promise as ideal
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candidates, such as graphene (Gr) [10,11], hexagonal boron nitride ( h-BN) [12-14] and MXenes
(derivatives of layered ternary carbides and nitrides known as MAX phases) [15,16]. They not
only have outstanding intrinsic properties, but also are geometrically compatible with lamella in
the “brick-and-mortar” structure. Meanwhile, these 2D nanomaterials exhibit high in-plane
rigidity while large out-of-plane flexibility, and thus their strengthening efficiencies in
composites are strongly affected by the way the individual sheets are arranged. Therefore, in the
“brick-and-mortar” structure, 2D nanomaterials are parallel arranged and their strengthening
capability could be fully exerted when a load is applied along the direction of their maximum
performance [17].
Among the aforementioned 2D nanomaterials, graphene is the most investigated emerging
nano-reinforcement for composites. Up to now, most of the studies on Gr/metal composite emphasize the homogeneous dispersion of grapheneMANUSCRIPT and the suppression of its restacking, but the orientation of graphene is random and out of control, such as in the Gr/metal composites
fabricated through ball-milling [18-22], molecular-level mixing [23,24], slurry blending [25,26] and electrodeposition [27]. The major difficulties for dispersing graphene in a metal matrix lie in graphene restacking, structural damage as well as poor interfacial bonding. Because of its 2D atomic layer structure, graphene exhibits high surface energy, high aspect ratio and strong Van der Waals interaction, and is prone to agglomerate during its mixing with metal [22,28].
Therefore, graphene is usually dispersed in metal matrix by using a ball-milling method. However, high energyACCEPTED during ball-milling will lead to structural damage of graphene [29] and some adverse reactions at interface [30]. Moreover, wettability of graphene with metal has to be improved to enhance interfacial bonding strength [31]. Beyond above difficulties, for fabricating
Gr/metal composites with the “brick-and-mortar” structure, additional bigger challenge is
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alignment of graphene in metal matrix, because graphene is flexible and has large aspect ratio.
Recently, efforts have been made to fabricate metal matrix composites reinforced with aligned graphene (or reduced graphene oxide, rGO), demonstrating the advantage from the resulting nanolaminated structure in strengthening and toughening the composites, but these fabrication processes are relatively complicated and difficult to be scaled-up [32-34]. To fabricate bulk
Gr/metal nanolaminated composites for large-scale applications, flake powder metallurgy has been established in our group in the past few years [35-39]. It is a bottom-up assembly process of composite flaky powders, i.e. nanoflake metal powders covered with rGO were used as building blocks to be orderly assembled together forming composites. This strategy has made a great success in strengthening and toughening lightweight Al matrix composites. Nevertheless, some major issues are not fully understood and addressed yet. Firstly, it remains a challenge to evaluate the rGO’s individual contribution in strenMANUSCRIPTgthening and toughening composite, because Al is chemically active and Al 2O3 passivation layer inevitably forms during preparing Al nanoflakes via a ball-milling process [35,40,36]. Secondly, although catalytically grown graphene has better intrinsic mechanical and functional properties than graphene oxide (GO, usually produced by chemical methods) [41-43], it is GO that is used more often in reinforcing
metal matrix because of its lower fabrication cost and easy handling. Therefore, a strategy is
desired to incorporate catalytically grown graphene in such bioinspired nanolaminated
composites for further improving overall performance. Thirdly, in contrast to intensive investigations onACCEPTED mechanical properties, functional properties in such bioinspired composites remain largely unexplored, while a synergy between them is required for more and more
emerging applications.
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Motivated by the aforementioned major issues, nacre-inspired Cu matrix nanolaminated composites reinforced with in-situ catalytically grown graphene were prepared in this study. Cu matrix composite has been intensively investigated for both structural and functional purposes
[44,45]. Unlike chemically active Al, copper oxide impurities are easy to be reduced. Moreover,
Cu is a typical catalyst for growing graphene, and Cu in various forms can be used for this purpose, such as foil [46], nanowire [47], particle [48-51] and porous Cu [52,53]. Here, Cu
submicron flakes cladded with in-situ catalytically grown graphene were used as the building
block for fabricating bulk Gr/Cu composite with the bioinspired nanolaminated structure via a
bottom-up assembly process. Tensile test on the as-obtained composites reveals that graphene in
the nanolaminated composites shows remarkably higher strengthening and stiffening efficiencies
than those of other reinforcements, while the composites retains a ductility and electrical
conductivity comparable to that of pure Cu. The superior overall performance is interpreted in terms of architecture effect, interfacial bonding and MANUSCRIPT interaction between graphene and Cu matrix. This work highlights the importance of architecture design in developing composites with good
combined structural and multifunctional properties.
2. Experimental methods
2.1. Fabrication of Cu flaky powders
100 g of commercially available spherical Cu powders (99.99% purity, with an averaged particle size of 40ACCEPTED m) were ball milled in a stainless steel mixing jar at a speed of 423 rpm for 5 h in pure ethanol, with a mass ratio between the Cu powder and the stainless steel milling ball of about 1:20.
2.2. In-situ fabrication of graphene/Cu flaky composite powders
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30g as-prepared Cu flaky powders were mixed with 250 ml 0.05~0.5 wt% Poly(methyl methacrylate) (PMMA) (M.W. 35000,) anisole solution. The slurry was stirred for 12 h, and then centrifuged at 4000 rpm for 10 min. The obtained PMMA/Cu flaky powders were dried in a vacuum oven at 85 ºC for 2 h to remove the solvent. The PMMA coated Cu flaky powders were put in an aluminum crucible and were heated in a tube furnace. Temperature was rapidly elevated to 900ºC under H 2 (100 sccm) and Ar (400 sccm) flow at atmospheric pressure and then kept at this temperature for one hour. The Gr/Cu composite powders were obtained by fast-cooling to room temperature under the H 2/Ar atmosphere.
2.3. Characterization of Gr/Cu flaky composite powders
Graphene was detached from the composite powder by etching the Cu substrate in a ferric chloride aqueous solution, then filtered and washed by water and ethanol, and finally dried in 60 ºC. Raman spectroscopy (Bruker Optics Senterra R200 MANUSCRIPT-L) was performed by using Ar + laser with a wavelength of 532 nm as excitation source to characterize structural integrity of graphene in composite powder. X-ray photoelectron spectroscopy (XPS, Kratos AXIS UltraDLD) was used to characterize elemental composition and chemical bonding. The morphology and distribution of graphene on the surface of Cu flaky powder were characterized by scanning electron microscope (SEM, Hitachi S-4800).
2.4. Consolidation of Gr/Cu composite powders
The Gr/Cu flakyACCEPTED composite powders were assembled in a graphite die (Ø30 mm) and then compacted in hot-pressing furnace. The powders were sintered in Ar atmosphere at 900 ºC and
50 MPa for 20 min under a heating rate of 15 ºC/min. The hot-pressed Gr/Cu composites were then hot-rolled with a reduction of 70% at 850 ºC for mechanical and electrical characterizations.
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2.5. Characterization of nacre-inspired Gr/Cu nanolaminated composite
Microstructure characterization of the Gr/Cu nanolaminated composite was performed using
SEM and HR-TEM, (JEOL JEM-2100F). The tensile testing was carried out on universal testing machine (Zwick/Roell Z100). Electrical conductivity was tested by four point probe instrument
(Ecopia EPS-300), and all the samples were wire-cut into a dimension of 10 mm × 5 mm × 0.2 mm and polished by 0.5- m Al 2O3 sandpaper to avoid the influence of rough surface on
conductivity measurement. Roughness of fractured surface was characterized by 3D optical
surface profile (ZeGage 3D Optical Surface Profiler).
2.6. Synthesis of Cu/Gr/Cu model materials and its interfacial tensile test
The model material was made of two pieces of copper foil and graphene in the middle.
Graphene was introduced on the surface of a 45- m thick Cu foil (20 mm × 10 mm) by using a MANUSCRIPT chemical vapor deposition (CVD) method, and the coverage was over 95%. The sandwiched Cu- foil/Gr/Cu-foil was obtained by sintering (700ºC and 50 MPa) the CVD Gr/Cu foil with one
piece of annealed Cu foil (suffered the same temperature condition in the CVD process). A
controlled sample of Cu-foil/GO/Cu-foil was prepared by sintering a GO covered Cu foil with
another annealed Cu foil. GO aqueous solution was spread on the surface of annealed Cu foil
and then dried. 0.05 wt% GO aqueous solution was used and adsorption time was optimized to
control the coverage of one-layer graphene oxide over 95%. Interfacial tensile tests on the model composites wereACCEPTED carried out on universal testing machine at a constant rate of 1.0 mm/min at room temperature until the two pieces of Cu foil was completely separated.
3. Results and discussion
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3.1. Synthesis and Microstructural Characterization
The fabrication process used to make nacre-inspired Gr/Cu composites is illustrated in Fig. 1.
Fig. 1. Schematic illustration of fabrication of Gr/Cu composite with nacre-inspired structure. Spherical Cu powder (a) were first transformed into MANUSCRIPT Cu flake (b) by a ball-milling process. (c) The as-obtained Cu flakes were soaked in an anisole solution of PMMA (typically less than
1wt %) and then dried in vacuum, forming a uniform PMMA film on the surface. (d) The coated
PMMA was used as carbon source for in-situ growing graphene at elevated temperature. (e) The
Gr/Cu composite powders were self-assembled into green compact by gravity because of its
large aspect ratio. (f) A nacre-inspired composite was finally obtained by a hot-pressing and hot-
rolling process.
First, starting fromACCEPTED spherical Cu powder ( Fig. 1 a), Cu flakes ( Fig. 1 b) were obtained by using a ball-milling method. The thickness of Cu flakes could be controlled by adjusting ball-milling
time and rotating speed. PMMA was coated on the surface of Cu flakes ( Fig. 1 c) and used as the
carbon source for graphene growth. Cu is a typical catalyst for growing graphene, and it acts as
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both catalyst and matrix in this study. By heating the PMMA coated Cu flakes at elevated temperature in a tube furnace under H 2/Ar flow (20 vol % H 2), graphene was grown on the surface to form Gr/Cu composite flakes (Fig. 1 d). The layer number of grown graphene could be
controlled by adjusting the concentration of PMMA solution in the coating process, and few
layer graphene could be obtained with a concentration less than 1 wt%. Under uniaxial
compaction at room temperature, the as-obtained Gr/Cu composite flakes self-assembled into
green compact with an orderly laminated structure owing to their large aspect ratio ( Fig. 1e). A
vacuum filtration can promote the formation of laminated structure. The aspect ratio and
thickness of the flakes have a decisive role on the final architecture, and ball-milling parameters
have to be optimized for a nanolaminated composite. Finally, fully densified Gr/Cu bulk
composites with a bioinspired nanolaminated structure were produced via vacuum hot pressing
and hot rolling ( Fig. 1f). Details of the fabrication process and microstructure evolution can be found in Experimental section and Supporting Inform MANUSCRIPTation ( Fig. S1 ).
The as-obtained Gr/Cu composite powders were characterized by scanning electron
microscopy (SEM), Raman spectroscopy and X-ray photoelectron spectroscopy (XPS) (see Fig.
S2-4 in the Supporting Information). As observed by SEM, grain boundaries in the Cu flakes are
clearly seen beneath the as-grown graphene, and even twin boundaries can be distinguished in
graphene/Cu flakes prepared from low PMMA concentrations, indicating few-layer graphene
with high light transmission is obtained [54]. To explore the quality of in-situ grown graphene, graphene was detachedACCEPTED from the composite powder by etching Cu in a ferric chloride aqueous solution and characterized by Raman spectroscopy. The relative intensity between the D and G
peaks ( ID/IG) reflects the quality of graphene, and was measured to be 0.24 and 0.62 for two
representative flakes. The values are significantly smaller than that of most graphene oxide,
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indicating lower defect concentration in catalytically grown graphene. Furthermore, XPS scans confirmed the sp 2-hybridized carbon orbitals of graphene [55]. The bonding characteristics were analyzed based on high resolution XPS, indicating no chemical bonding between graphene and
Cu in the as-obtained
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Fig. 2. Structure of nacre-inspired Gr/Cu composite compared to natural nacre. Panels (a–c) correspond to nacre;ACCEPTED panels (d –f) to nacre-inspired Gr/Cu composite. The as-obtained Gr/Cu composites show a similar “brick-and-mortar” structure as that in natural nacre. (a,d) SEM
micrographs showing the long-range order of flakes, and (b,e) local stacking of flakes. The insets
in (a,d) are nacre and densified Gr/Cu composite, respectively. (c,f) TEM micrographs showing
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local stacking of flakes. (g) An EBSD image of the cross-section of the nacre-inspired Gr/Cu composite, revealing a laminated structure, from which boundary spacings of the elongated Cu grains were estimated. (h,i) Distribution of the boundary spacing parallel ( dT) and perpendicular
(dL) to the hot-pressing direction. At least 150 boundaries were measured and statistically averaged.
Gr/Cu composite powders. Weak interaction between graphene and Cu in an as grown CVD
Gr/Cu foil is confirmed by a low adhesion energy value of ~0.72 J m −2 obtained from a double
cantilever beam (DCB) test on the interface [56]. The measured value is close to other such
values calculated through quantum simulations according to a van der Waals interaction [57,58].
Fig. 2 highlights strong similarities at several length scales between natural nacre ( Fig. 2 a-c) and the Gr/Cu composites after a hot-processing process ( Fig. 2 d-f), validating the possibility to fabricate bulk, centimeter-sized or even larger sam plesMANUSCRIPT with only a few simple processing steps. In natural nacre, ~500 nm thick platelets of mineral aragonite are compared to “bricks”, bonded by ~10 nm thick protein mortar in-between, resulting in a regular “brick-and-mortar” structure.
The final microstructure of nacre-inspired Gr/Cu composite is characterized by a dense stacking of lamellae presenting long-range order, in which 2D graphene and Cu flake are alternately stacked. Each Cu matrix lamella contains predominantly a single grain through its thickness. As revealed by the electron backscatter diffraction (EBSD) analysis, no obvious crystallographic texture exists in hot-pressed composites ( Fig. 2g). The distribution of boundary spacing scaled
by the interceptionACCEPTED length along lines perpendicular ( dT) and parallel ( dL) to the hot-pressing direction is shown in Fig. 2h and Fig. 2i, respectively. The average sizes of the elongated grains are dT ≈ 0.66 and dL ≈ 1.66 m. Without the confinement by graphene, the grains in the
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unreinforced Cu matrix recrystallized and were found to be equiaxialed crystals with an average size of ~2.02 m (see Fig. S5 in the Supporting Information).
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Fig. 3. Tensile properties and electrical conductivity of Gr/Cu composites. (a) Engineering stress-
strain curves forACCEPTED Gr/Cu composites and the unreinforced Cu matrix. (b) Comparison of the
strengthening and stiffening efficiencies of graphene in Gr/Cu composites with various
reinforcements in other reported Cu matrix composites. (c) Comparison of overall properties
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including tensile strength, elongation and electrical conductivity of the Gr/Cu composites with other reported Cu matrix composites. The curved shadows at the bottom indicate the conflict between strength and elongation.
Table 1 . Tensile properties and electrical conductivity of the Gr/Cu nanolaminated composites
and the unreinforced Cu matrix.
sample yield tensile Young’s total electrical strength strength modulus elongation conductivity (MPa) (MPa) (GPa) (%) (%IACS) a Cu 72 218 108 43.5 97.8 1.6 vol %- 122 305 127 41.0 97.1 Gr/Cu 2.5 vol %- 200 378 135 32.3 93.8 Gr/Cu
aIACS is an acronym for international annealed copper standard, 58 × 10 6 S m ‒1 at 20 ºC.
MANUSCRIPT 3.2. Mechanical and Electrical Properties
Tensile tests were performed to evaluate the strength, stiffness, and elongation of the rolled
Gr/Cu composites. Fig. 3 a shows representative engineering stress-strain curves of the nacre-
inspired 1.6 and 2.5 vol % Gr/Cu nanolaminated composites (the volume fractions of graphene
were estimated by measuring average thickness of graphene layer and Cu matrix layer in
transmission electron microscope (TEM)), together with that of unreinforced Cu matrix fabricated using ACCEPTEDidentical processing conditions but without graphene growth. Key mechanical data obtained from the tensile tests are summarized and tabulated in Table 1 . 1.6 vol % Gr/Cu
composite was shown to have a tensile strength of 305 ± 10 MPa and a Young’s modulus of 127
± 3 GPa, ~40% and ~18% higher than those of the Cu matrix, respectively. Tensile strength and
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Young’s modulus of the composite further increase to 378 ± 8 MPa and a Young’s modulus of
135 ± 4 GPa (~73% and ~25% enhancement over the Cu matrix) when the graphene concentration increases to 2.5 vol %. These results clearly indicate that graphene is a highly effective reinforcement in the nacre-inspired nanolaminated composites. The high strengthening capability can be better presented by comparing its strengthening and stiffening efficiencies with those of Cu matrix composites reinforced by other reinforcements. The strengthening (or stiffening) efficiency of a reinforcement in a composite is defined as the strength (or modulus) increment per unit volume fraction of the reinforcement, i.e ., (σc – σm)/ Vfσm and ( Ec – Em)/ VfEm,
where σc and σm are the tensile strengths of the composite and the matrix, respectively, Ec and Em are the Young’s modulus of the composite and the matrix, respectively, and Vf is the volume fraction of the reinforcement. As shown in Fig. 3b, the strengthening and stiffening efficiencies
of graphene in the nacre-inspired nanolaminated composites are considerably higher than those of conventional particle and fiber reinforcements, andMANUSCRIPT also superior to those of carbon nanotube (CNT) as well as graphene (or its derivatives) in other Cu matrix composites reported so far [59-
69].
The simultaneous attainment of both high strength, high ductility as well as high conductivity
in metals is a vital requirement for many modern applications; unfortunately, methods used to strengthen metals generally also cause a pronounced decrease in ductility and electrical conductivity. In this work, the tensile tests show that the uniform elongation of 41 ± 1.5% for the 1.6 vol% Gr/Cu ACCEPTEDand 32.3 ± 1.6% for 2.5 vol% Gr/Cu nacre-inspired nanolaminated composites are moderately lower than that of pure Cu matrix (43.5 ± 0.7%). At the same time, the electrical conductivity of the composites is identical to that of the Cu matrix. As shown in Fig. 3c, a good balance between strength, ductility and conductivity has been achieved in the in-situ grown
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Gr/Cu nacre-inspired nanolaminated composites [70-80].
3.3. Interface Structure and Bonding
Fig. 4a is a representative TEM image for the interface between graphene and the copper
matrix, showing that the interfaces are free of impurities, voids, or gaps. The few-layer graphene
(11 layers) shows a distinct lattice fringe with a 0.34 nm spacing of (0002) plane, confirming
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Fig. 4. Analyses on Gr/Cu interfaces. (a) TEM micrograph image showing the Gr/Cu interface.
(b) High resolution TEM image of Gr/Cu interface along the [110] direction of Cu matrix. The
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corresponding FFT is shown in the upper right inset. (c) GPA map (local g-map) of the (b) showing the position of the misfit dislocations at the interface as hot spots. (d) Inverse FFT of one family of planes to show the position of dislocations indicated as extra half planes (pointed out by “T” symbols). Both local g-map (c) and inverse FFT (d) were obtained using the g =
{111} indicated by circle in the FFT inset of (b). they are highly graphitized and of high quality. The thickness of graphene could be controlled by adjusting the carbon source concentration of PMMA solution, but obtaining graphene with less than five layers is still of challenge by using the solid carbon source (see Fig. S6 in the
Supporting Information). According to full atomistic nanoindentation simulations, Chang et al
[81] reported that the hardness of the Ni/Gr/Ni sandwiched nanocomposites decreases with increasing the numbers of graphene layers ( N). The weak Van der Waals interaction between graphene sheets is responsible for the reduction. InMANUSCRIPT our cases, however, increasing N of in-situ grown graphene increases mechanical strength. The possible reason for the opposite conclusions may be the discrepancy between the volume fractions of graphene in the two works. In Chang’s
Ni/Gr/Ni sandwiched nanocomposites, the volume fraction of graphene ranges from 4.3 to 21 vol % (the number of graphene layer varies from N = 1 to N = 5 while the thickness of metal layer was kept constant as 8 nm), which is one order of magnitude higher than our samples (such as 1.6 and 2.5 vol %). When volume fraction of graphene is as low as that in this work, the strength increment contributed by increasing volume fraction can compensate the loss caused by the weak Van derACCEPTED Waals interaction between graphene nanosheets. However, for a fixed volume fraction of graphene (both high and low cases), Chang’s conclusion suggests that graphene with fewer layers is beneficial for enhancing mechanical strength, because it means refinement of metal matrix layer (namely grain size) and reduction of adverse effect of weak Van der Waals
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forces. Therefore, preparing a composite with thinner metal matrix layer thickness ( λ) and less
graphene layer number ( N) is suggested [32].
As revealed by geometric phase analysis (GPA) map and inverse fast Fourier transform (FFT)
image, misfit dislocations along the Gr/Cu interface could be easily visible. GPA is an image
processing technique which is sensitive to small displacements of the lattice fringes in high-
resolution TEM (HR-TEM) images and dislocations are shown as hot-spots [82]. Fig. 4b shows a
typical HR-TEM image of the Gr/Cu interface along the [110] zone axis of Cu matrix in the as-
obtained nacre-inspired composites. The corresponding FFT was also shown in upper right inset
of Fig. 4b. Fig. 4c shows the corresponding GPA evaluation on Fig. 4b. The color map
represents the strain component according to the εxy [111] direction. Most part of the graphene shows uniform strain and no strain difference, which is chosen as the reference frame. The matrix has a color between red and green, corresponMANUSCRIPTding to about –16% lattice strain when compared with the reference, which is too large for a lattice accommodation via elastic strain.
Indeed, as seen in the corresponding inverse FFT image ( Fig. 4d) of one family of planes (shown by a circle in the FFT inset in Fig. 4b), periodic interface mismatch dislocations can be clearly recognized as extra half planes which are indicated by “T” symbols. Such lattice mismatch dislocations between the graphene and the Cu matrix occur every five or six {111} planes, namely 16~20% mismatch, corresponding to the ~16% lattice strain estimated from the color map ( Fig. 4 c). The misfit dislocation with high density at interface plays an important role in strengthening nacre-inspiredACCEPTED Gr/Cu composites because of stronger interaction between dislocations during metal deformation.
Interface bonding is of great importance in composites because it affects the load-transfer and
energy-exchange efficiency between reinforcement and matrix. As far as we know, graphene and
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its derivatives have been used as reinforcements in metal matrix composites, but no direct experimental data has been reported to compare their relative interface bonding strength in composites. Conventional methods such as push -out and pull -out are effective for measuring the interface bonding strength of traditional ceramic and carbon micro -fiber reinforcement in composites, but they are of great challenge [83] or even impossible for flexible nano
Fig. 5. (a) Schematic illustration of fabrication processMANUSCRIPT of Cu -foil/Graphene/ Cu -foil and Cu - foil/Graphene oxide/Cu -foil specimens for interface tensile testing. (b) Comparison on interface bonding strength in the composites reinforced with CVD graphene and adsorped graphene oxide. reinforcements such as carbon nanotube and graphene. Recently, Hwang et al reported the direct measurement of adhesion energy between CVD graphene and Cu in a Cu -foil/Gr/Cu -foil sandwiched model composite by using a DCB fracture test [23].
Here, to compare the relative value of interface bonding strength between catalytically grown graphene and grapheneACCEPTED oxide, a simplified interface tensile testing based on DCB was carried out on two model composites of Cu -foil/CVD Gr/Cu -foil and Cu -foil/GO/Cu -foil. Fig. 5a shows the fabrication process for the two model composites. Monolayer of both CVD graphene and graphene oxide were used to ensure that the interface bonding strength between the graphene and
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Cu, but not between graphene layers, was measured. The graphene coverage rate on Cu foil in two model composites was also controlled to be commensurate with each other and above 95% for comparable fraction of graphene/Cu interface. Although no precise value of adhesion energy as in the standard DCB fracture test was measured, relative interface bonding strength was compared in these two model composites as indicated by force-displacement curves obtained under the same testing conditions ( Fig. 5 b). The energy required to separate the bonded Cu foils can be roughly estimated based on the area under the force-displacement curves, and the result shows that the energy for the interface with CVD graphene is about 80% higher than that for the interface with graphene oxide. The difference on interface bonding strength can be also appreciated by comparing the fractured surfaces (see Fig. S7 in the Supporting Information). The
fractured interface with CVD graphene is very rough and shows many ravines, which means
interface bonding is strong enough to cause failure of Cu matrix during detachment, while the fractured interface with graphene oxide is smoother MANUSCRIPT than the former. The adhesion energy between graphene and Cu as grown from CVD can be enhanced after sintering at high
temperature and high pressure, which was explained by the existence of native oxygen on Cu and
the formation of strong oxygen mediated carbon-Cu covalent bonding [23]. Along this line of
consideration, the bonding strength in the interface with graphene oxide would be stronger than
that with CVD graphene because of higher oxygen content, but the conclusion is reverse in our
measured data. A possible reason might be the difference between lateral sizes of CVD graphene
and graphene oxide. The lateral size of CVD graphene is about 10~20 m, an order of magnitude larger than that ofACCEPTED the graphene oxide prepared by the Hummers’ method. It need more energy to tear a continuous CVD graphene than fragmented graphene oxide. The difference on lateral size makes the comparison complicated, but it implies that larger lateral size of graphene can
20 ACCEPTED MANUSCRIPT
compensate the loss of interface bonding strength owing to the absence of bridge oxygen as that in graphene oxide. Moreover, the catalytically grown graphene has better intrinsic mechanical and functional properties than graphene oxide.
3.4. Fracture Behavior
Fig. 6 shows fracture surface for the nacre-inspired Gr/Cu composites and pure Cu matrix.
The pure Cu sample shows well-developed dimples over the entire fractured surface, indicating a
typical ductile fracture with high plastic deformation. No laminated structure was found in the
pure Cu sample because of grain growth during consolidation at high temperature. With
introducing graphene and increasing its volume fraction, the nacre-inspired nanolaminated
structure was formed. As we can see from the fracture surface morphology, there is no graphene
pull-out on the fractured surface, which is different from other reported graphene oxide reinforced metals [33]. The different facture behaviorMANUSCRIPT could be understood by a simple shear lag model [84]. The failure behavior (either pull-out or fracture of reinforcement) is determined by
its relative size compared to the ratio of tensile strength of reinforcement to the yield shear
strength of matrix. In this work, the yield shear strength of Cu matrix should have no substantial
difference with that in other reported composites, and the difference might originate from
competition between the strength and size of graphene. Compared to the reduced graphene oxide
used in most reported metal matrix composites, the graphene in this work has higher structural
integrity and therefore higher tensile strength, which means a larger critical aspect ratio in the Gr/Cu composites.ACCEPTED On the other hand, as estimated from the surface morphology analysis on the Gr/Cu flaky powder from SEM images, the lateral size of the graphene here is on the scale of
several micrometers at least. This size is remarkably larger than that of the graphene oxide used
in most reported composites, resulting in a predominant failure behavior of graphene fracture but
21 ACCEPTED MANUSCRIPT
not pulling-out.
MANUSCRIPT Fig. 6. Fractured surface analysis on (a-c) unreinforced copper matrix, (d-f) 1.6 vol % Gr/Cu,
and (g-i) 2.5 vol % Gr/Cu by using 3D optical surface profiler (the left column) and SEM with
low magnification (the middle column) and high magnification (the right column). The value of
Sa in the left column is a 3D roughness parameter based on arithmetic mean height in the
measured area.
3D optical surface profile was used to compare the fractured surface roughness in three samples. As indicatedACCEPTED by the value of Sa (see Fig. S8 in the Supporting Information), calculated from the arithmetic mean amplitude of fractured surface fluctuation, substantial discrepancy on surface roughness was measured between the three samples ( Fig. 6a, d, g). As compared to unreinforced copper matrix, the nacre-inspired nanolaminated architecture caused an increase on
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surface roughness by 108% and 161% in the 1.6 vol % Gr/Cu and 2.5 vol % Gr/Cu, respectively, indicating that the bioinspired architecture increased the total fracture surface area in the composites and therefore resulted in greater energy absorption as compared to an unreinforced matrix.
3.5. Discussions
The advantage of the nacre-inspired nanolaminated structure for balancing mechanical properties and electrical conductivity lies in both architecture effect and improved interface bonding ( Fig. 7). The nacre-inspired nanolaminated architecture provides extrinsic toughening by crack deflection as indicated by increased fractured surface area. Meanwhile, with this
MANUSCRIPT
Fig. 7. Structure-function integration and synergy benefit from bioinspired nanolaminated architecture coupledACCEPTED with improved interface bonding in Cu matrix composite reinforced with in- situ grown graphene. Nanolaminated architecture not only toughen the composite by deflecting
crack propagation, but also align graphene to maximize its performance for required loading and
carrier transporting conditions. The in-situ catalytic growth process provides high structural
23 ACCEPTED MANUSCRIPT
quality and enhanced interface bonding strength after sintering. nanolaminated architecture, aligning graphene along the direction for required loading and carrier transporting conditions facilitates the load transfer between graphene and Cu matrix as well as enhancement on electrical conductivity. At the same time, the presence of high surface area graphene increases thermal stability of the nanolaminated structure and retains it during its hot-pressing and rolling at high temperature, while the grains coarsen in the pure Cu without graphene. The overall properties are expected to be further enhanced by optimizing the geometrical parameters, such as the layer number of grown graphene and the thickness of Cu matrix lamella. The former could be adjusted by using a gas carbon source, while the latter controlled by varying the conditions of ball-milling and deformation processing.
Another characteristic is the multipole role of Cu, acting as both the catalyst of growing graphene and metal matrix in the composites. Cataly MANUSCRIPTtically grown graphene is desired because higher structural quality as well as improved interface bonding. In general, the wettability of pure
Cu on carbon is poor [85], making it difficult to disperse graphene in Cu and to form strong bonding with Cu matrix via casting or ball-milling processes, which could be resolved by in-situ growing graphene on Cu flakes. Moreover, this bioinspired strategy provides an idea for alignment and assembly of graphene in metal matrix. Also, this strategy could be extended to other metal catalysts that can be used for growing graphene, such as Fe, Co, Ni and their alloys as well as some noble metals, and scaled up for mass production because it is compatible with conventional powderACCEPTED metallurgy technology.
4. Conclusions
In this work, we have fabricated in-situ catalytically grown Gr/Cu composites with nacre-
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inspired nanolaminated architecture. The resulting composites show a synergy between mechanical strength, elongation, elastic stiffness and electrical conductivity. The bioinspired nanolaminated architecture toughens the composite by crack deflection, strengthens the composites and retains electrical conductivity by alignment of 2D catalytically grown graphene to maximize its performance for required loading and carrier transporting conditions.
Meanwhile, in-situ grown graphene improves interface bonding and structural quality. The strategy could be extended to other metal matrixes and multifunctional materials for thermal conductivity, tribological and electrical contact applications, shedding light on developing metal matrix composites with good combined structural and multifunctional properties.
Acknowledgments
The authors would like to acknowledge the Natural Science Foundation of China (Nos.51371115, 51671130, 51131004), the Ministry of ScienceMANUSCRIPT & Technology of China (973 program, No.2012CB619600), Shanghai Science & Technology Committee (Nos.14JC1403300,
14DZ2261204, 15JC1402100, 14520710100).
Appendix A. Supplementary data Supplementary data related to this article can be found at ………..
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